A modern, advanced design industrial gas turbine (IGT) has hot stage blades and vanes which are required to perform for lives of 2 to 5.times.10.sup.4 up to 10.sup.5 hours, e.g. at least about 30,000 hours in a corroding environment resulting from the combustion of relatively low grade fuels and, in the case of blades, under high stress. Naturally, in order to increase efficiency, it is desired to operate such IGT blades and vanes at the higher practical operating temperatures consistent with achieving the design life-times. When considering operating temperatures, it is necessary to take into account not only the highest temperature to which a turbine blade is exposed, but also a range of temperatures below that highest temperature. Even at steady-state operation, a turbine blade will experience a variety of temperatures along its length from root to tip and across its width from leading to trailing edge.
Over the long design lives of IGT blades and vanes, corrosion resistance and oxidation resistance become more important factors than they are in the well-developed field of aircraft gas turbine (AGT) alloys. Although in neither the case of AGT nor IGT turbine blades or vanes would it be advisable to select an oxidation or corrosion prone alloy, the longer (by an order of magnitude) time exposure of IGT components to a more corroding atmosphere make oxidation and corrosion resistance very important features of IGT alloy structures. IGT alloy structures such as hot stage blades and vanes can be coated with conventional coatings to enhance oxidation and corrosion resistance but these coatings are subject to cracking, spalling and the like. Over the long design lives of IGT components, it is more likely that coating failures will occur in comparison to such failures with AGT coated components used for shorter time periods. Thus, even if coated, an IGT alloy structure used in the hot stage of an IGT must have the best oxidation and corrosion resistance obtainable commensurate with other required properties and characteristics.
In designing alloy structures for IGT turbine blades it is natural to investigate nickel-base alloys which are used conventionally in AGT turbine blades. Even the strongest conventional, .gamma.' strengthened nickel base alloys rapidly lose strength at temperatures above about 900.degree. C. (see FIG. 2 of U.S. Pat. No. 4,386,976). It is disclosed in U.S. Pat. No. 4,386,976 however that nickel-base alloys combining .gamma.' strengthening and strengthening by a uniform dispersion of microfine refractory oxidic particles can provide adequate mechanical properties in the temperature range of 750.degree. C. up to 1100.degree. . However, the alloys disclosed in U.S. Pat. No. 4,386,976 are deemed to have inadequate oxidation and corrosion resistance for use in advanced design IGTs. It is also known, for example, from U.S. Pat. No. 4,039,330 that .gamma.' strengthened nickel-base alloys containing in the vicinity of 21 to 24 weight percent chromium along with some aluminum have excellent corrosion and oxidation resistance, of the character needed for IGT usage. At very high temperatures, e.g. over 1000.degree. C., the oxidation resistance of alloys as disclosed in U.S. Pat. No. 4,039,330 tends to fall off. Strength at temperatures in excess of 900.degree. C. of the alloys disclosed in U.S. Pat. No. 4,039,330, as with all .gamma.' strengthened nickel-base alloys is inadequate for components of advanced design IGTs.
From the background in the immediately preceding paragraph one might be tempted to declare that the solution to providing turbine blades for advanced design IGTs is obvious. Either increase the chromium and/or aluminum content of .gamma.' and dispersion strengthened alloys disclosed in U.S. Pat. No. 4,386,976 or add dispersion strengthening to the alloys disclosed in U.S. Pat. No. 4,039,330. These appealing, seemingly logical solutions to the existing problem are overly simplistic.
The first possibility i.e., increasing the chromium and/or the aluminum content of a known .gamma.' and dispersion strengthened alloy, has two difficulties. Increasing either chromium or aluminum can tend to make a nickel-base alloy sigma prone. Increase of chromium directly dilutes the nickel content of the alloy matrix remaining after .gamma.' phase precipitation. Increasing the aluminum content increases the amount of .gamma.' phase (Ni.sub.3 Al-Ti) which can form in the nickel-base alloy again diluting the matrix with respect to nickel. Detrimental acicular sigma phase tends to form in nickel-base alloys having low nickel matrix contents after intermediate temperature (e.g., 800.degree. C.) exposure resulting in low alloy ductility. Because the existence of .gamma.' phase is essential to component strength at temperatures up to about 900.degree. C., it is necessary to carefully control alloy modification to avoid phase instability over the long term usage characteristic of IGTs where a minimum acceptable ductility is essential. From another point of view, indiscriminate alloy modification especially in the realm of increasing aluminum and/or chromium contents presents a difficulty in providing the component microstructure essential to strength of dispersion strengthened alloys at high temperature. Referring again to U.S. Pat. No. 4,386,976 Column 1, line 58 et seq., it is disclosed that ODS (oxide dispersion strengthened) alloys must be capable of developing a coarse, elongated grain structure in order to obtain good elevated temperature properties therein. This coarse, elongated grain structure is developed by directional, secondary recrystallization at a temperature above the .gamma.' solvus temperature and below the incipient melting temperature of the alloy (see Column 6, line 58 et seq. of the U.S. Pat. No. 4,386,976) or some temperature close to the incipient melting temperature. If .gamma.' phase is not solutioned, the secondary crystallization will not proceed. If the incipient melting temperature of the alloy is exceeded the oxide dispersion will be detrimentally affected. For practical production, the interval between the .gamma.' solvus temperature and the temperature of incipient melting must be at least about 10.degree. and, more advantageously, at least about 20.degree. in celsius units. Because of the complexity of modern .gamma.' strengthened alloy compositions and the complex interactions among the alloying elements, there is no way of predicting the secondary recrystallization interval which is a sine qua non for obtaining the high temperature strength in ODS alloys.
The same difficulty applies to the possible idea of providing oxide dispersion strengthening to a known, high strength .gamma.' oxidation and corrosion-resistant alloy. There is no way of predicting whether nor not the theoretical ODS-.gamma.' strengthened alloy can be made on a commercial basis.
The foregoing makes it clear that the provision of alloy components suitable for hot stage advanced design IGT usage is a problem that requires critical metallurgical balancing to at least provide an adequate window for thermal treatment necessary for practical production of such components. In addition, the alloy composition must be capable of undergoing the practical mechanical and thermomechanical processing required to reach the stage of directional recrystallization.
The present invention provides alloy bodies suitable for use in advance design IGTs which can be produced in a practical manner.